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Materials <strong>and</strong> Design 43 (2013) 373–383<br />
Contents lists available at SciVerse ScienceDirect<br />
Materials <strong>and</strong> Design<br />
journal homepage: www.elsevier.com/locate/matdes<br />
<strong>Tensile</strong> <strong>and</strong> <strong>fatigue</strong> <strong>properties</strong> <strong>of</strong> <strong>fiber</strong> <strong>laser</strong> <strong>welded</strong> <strong>high</strong> <strong>strength</strong> <strong>low</strong> alloy<br />
<strong>and</strong> DP980 dual-phase steel joints<br />
W. Xu a , D. Westerbaan b , S.S. Nayak b , D.L. Chen a,⇑ , F. Goodwin c , Y. Zhou b<br />
a Department <strong>of</strong> Mechanical <strong>and</strong> Industrial Engineering, Ryerson University, 350 Victoria Street, Toronto, Ontario M5B 2K3, Canada<br />
b Department <strong>of</strong> Mechanical <strong>and</strong> Mechatronics Engineering, University <strong>of</strong> Waterloo, 200 University Avenue West, Waterloo, Ontario N2L 3G1, Canada<br />
c International Zinc Association, Durham, NC 27713, USA<br />
article<br />
info<br />
abstract<br />
Article history:<br />
Received 2 June 2012<br />
Accepted 8 July 2012<br />
Available online 16 July 2012<br />
Keywords:<br />
Fiber <strong>laser</strong> welding<br />
High <strong>strength</strong> <strong>low</strong> alloy steel<br />
Dual-phase steel<br />
Microstructure<br />
Fatigue<br />
The study was aimed at evaluating the microstructure <strong>and</strong> mechanical <strong>properties</strong> <strong>of</strong> <strong>high</strong>-speed <strong>fiber</strong><br />
<strong>laser</strong> <strong>welded</strong> <strong>high</strong> <strong>strength</strong> <strong>low</strong> alloy (HSLA) <strong>and</strong> DP980 dual-phase steel joints with varying weld geometries.<br />
Fusion zone (FZ) consisted <strong>of</strong> martensitic structure, <strong>and</strong> heat-affected zone (HAZ) contained<br />
newly-formed martensite in both steels <strong>and</strong> partially tempered martensite in DP980. While HAZ-s<strong>of</strong>tening<br />
was present in DP980, it was absent in HSLA. A distinctive ‘‘suspension bridge’’-like hardness pr<strong>of</strong>ile<br />
with the FZ hardness as a ‘‘pylon’’ appeared in the <strong>fiber</strong> <strong>laser</strong> <strong>welded</strong> joints. Both HSLA <strong>and</strong> DP980 joints<br />
showed a superior tensile <strong>strength</strong>, with a joint efficiency <strong>of</strong> 94–96% <strong>and</strong> 96–97%, respectively, despite a<br />
reduced elongation in DP980 joints. Fatigue <strong>strength</strong> was <strong>high</strong>er in DP980 joints than in HSLA joints at<br />
<strong>high</strong>er stress amplitudes, but had no obvious difference at <strong>low</strong>er stress amplitudes. DP980 multiple linear<br />
welds exhibited a larger scatter <strong>and</strong> <strong>low</strong>er <strong>fatigue</strong> <strong>strength</strong>. Fatigue failure <strong>of</strong> HSLA joints occurred in the<br />
base metal at a stress amplitude above 250 MPa, <strong>and</strong> at weld concavity at a <strong>low</strong>er stress amplitude be<strong>low</strong><br />
250 MPa. Fatigue crack in DP980 joints initiated predominantly from the weld concavity at both <strong>high</strong> <strong>and</strong><br />
<strong>low</strong> levels <strong>of</strong> stress amplitudes.<br />
Ó 2012 Elsevier Ltd. All rights reserved.<br />
1. Introduction<br />
The dem<strong>and</strong> <strong>of</strong> environment-friendly vehicles that have better<br />
fuel economy <strong>and</strong> <strong>low</strong>er CO 2 emissions compels automakers to<br />
apply advanced technology <strong>and</strong> new materials in vehicle manufacturing<br />
[1–4]. During manufacturing <strong>of</strong> automotive components,<br />
welding <strong>and</strong> joining are unavoidable. Laser welding is an enabling<br />
technology which is fast <strong>and</strong> precise in joining a wide variety <strong>of</strong><br />
materials with varying thicknesses <strong>and</strong> types [5,6]. A new generation<br />
<strong>of</strong> <strong>fiber</strong> <strong>laser</strong>s has been developed for industrial applications<br />
<strong>and</strong> has multiple advantages, including <strong>high</strong> power with small<br />
beam divergence, flexible beam delivery, <strong>low</strong> maintenance costs,<br />
<strong>high</strong> efficiency, <strong>and</strong> compact size [7,8].<br />
A lot <strong>of</strong> studies on <strong>laser</strong> welding <strong>of</strong> HSLA, DP980, <strong>and</strong> other<br />
steels have been reported in the literature [9–12]. For example,<br />
Saunders <strong>and</strong> Wagone [9] conducted a study on Nd:YAG <strong>and</strong> CO 2<br />
<strong>laser</strong> welding <strong>of</strong> the aluminum-killed drawing quality (AKDQ) steel<br />
<strong>and</strong> HSLA steel sheets in which they concluded that the intrinsic<br />
ductility <strong>of</strong> AKDQ <strong>and</strong> HSLA steel reduced after <strong>laser</strong> welding. In<br />
another study, Sreenivasan et al. [10] investigated the mechanical<br />
<strong>properties</strong> <strong>and</strong> performance <strong>of</strong> DP980 steel tailor <strong>welded</strong> blanks<br />
manufactured using Nd:YAG <strong>and</strong> diode <strong>laser</strong> welding (DLW).<br />
⇑ Corresponding author. Tel.: +1 416 979 5000x6487; fax: +1 416 979 5265.<br />
E-mail address: dchen@ryerson.ca (D.L. Chen).<br />
Formability <strong>and</strong> uniaxial tensile <strong>strength</strong> <strong>of</strong> DP980 <strong>welded</strong> blanks<br />
significantly reduced due to the occurrence <strong>of</strong> heat affected zone<br />
(HAZ)-s<strong>of</strong>tening, i.e., the hardness at outer HAZ being significantly<br />
be<strong>low</strong> that <strong>of</strong> the base material caused by tempering <strong>of</strong> martensite.<br />
Similar effect <strong>of</strong> HAZ-s<strong>of</strong>tening on the formability was observed by<br />
Xia et al. [11] in their comparative study <strong>of</strong> HSLA <strong>and</strong> DP980 steel<br />
joints prepared using DLW. The study suggested a <strong>low</strong>er formability<br />
<strong>of</strong> DP980 diode <strong>laser</strong> <strong>welded</strong> blanks compared to the base metal<br />
(BM), which could be improved by increasing welding speed which<br />
in turn reduced the HAZ-s<strong>of</strong>tening. In addition, it was also reported<br />
that changes in the microstructure <strong>and</strong> hardness in the HAZ <strong>and</strong> FZ,<br />
in diode <strong>laser</strong> welds with similar <strong>and</strong> dissimilar combination <strong>of</strong><br />
DP600 <strong>and</strong> DP980 steel, resulted in a change <strong>of</strong> <strong>fatigue</strong> performance<br />
[13,14].<br />
The important outcome <strong>of</strong> the above mentioned studies was<br />
that microstructural heterogeneity (mainly HAZ-s<strong>of</strong>tening) significantly<br />
affects the mechanical <strong>properties</strong> such as tensile <strong>and</strong> <strong>fatigue</strong><br />
<strong>strength</strong> <strong>and</strong> also the formability <strong>of</strong> <strong>laser</strong> <strong>welded</strong> HSLA <strong>and</strong> DP980<br />
joints. High speed <strong>fiber</strong> <strong>laser</strong> welding (FLW) is becoming increasingly<br />
important <strong>and</strong> its share in the total <strong>laser</strong> industry is expected<br />
to grow substantially [8]. One <strong>of</strong> the reasons is that the <strong>high</strong> power<br />
density, small beam divergence <strong>and</strong> focal spot diameter could lead<br />
to welds with a <strong>high</strong>er penetration <strong>and</strong> smaller width for a given<br />
welding speed [7]. However, there is no report so far on the effect<br />
<strong>of</strong> FLW on the microstructure <strong>and</strong> mechanical <strong>properties</strong> <strong>of</strong> the<br />
0261-3069/$ - see front matter Ó 2012 Elsevier Ltd. All rights reserved.<br />
http://dx.doi.org/10.1016/j.matdes.2012.07.017
374 W. Xu et al. / Materials <strong>and</strong> Design 43 (2013) 373–383<br />
Table 1<br />
Compositions <strong>of</strong> the HSLA <strong>and</strong> DP980 steel used in the present study.<br />
Steel grade C Mn Si Al Mo Cr N<br />
HSLA 0.0795 0.827 0.454 0.048 0.007 0.033 0.007<br />
DP980 0.15 1.5 0.31 0.05 N/A N/A N/A<br />
important automotive steels namely HSLA <strong>and</strong> DP980. It is unclear<br />
how serious the HAZ s<strong>of</strong>tening in the HSLA <strong>and</strong> DP980 steels after<br />
FLW is, <strong>and</strong> if the multiple linear FLW would have a significant effect<br />
on the <strong>fatigue</strong> <strong>strength</strong>. This study was, therefore, aimed at<br />
identifying the evolution <strong>of</strong> microstructure <strong>and</strong> the effects <strong>of</strong><br />
FLW on the <strong>fatigue</strong> resistance in HSLA <strong>and</strong> DP980 steels. The effects<br />
<strong>of</strong> weld geometry, i.e., single <strong>and</strong> multiple linear welds, on the<br />
microstructure <strong>and</strong> <strong>fatigue</strong> <strong>properties</strong> <strong>of</strong> the <strong>welded</strong> joints have<br />
also been evaluated.<br />
2. Material <strong>and</strong> experimental procedure<br />
The starting materials in the present study were hot-dip galvanizing<br />
(GI) coated HSLA <strong>and</strong> DP980 steel with a thickness <strong>of</strong><br />
1.2 mm; the chemical compositions <strong>of</strong> the steels are given in Table<br />
1. Butt welding <strong>of</strong> the steel sheets was carried out using an<br />
IPG PhotonicsYLS-6000 <strong>fiber</strong> <strong>laser</strong> system with a power <strong>of</strong> 6 kW,<br />
a welding speed <strong>of</strong> 16 m/min, a focal length <strong>of</strong> 30 cm. The <strong>fiber</strong> <strong>laser</strong><br />
had a <strong>fiber</strong> core diameter <strong>of</strong> 0.3 mm with a <strong>laser</strong> beam spot<br />
size/diameter <strong>of</strong> 0.6 mm. The weld geometries <strong>and</strong> dimensions <strong>of</strong><br />
the steel sheets used in this study are shown in Fig. 1a. Welding<br />
was performed with 0° head angle, i.e., the <strong>laser</strong> beam was perpendicular<br />
to the surface <strong>of</strong> work pieces <strong>and</strong> without any shielding gas.<br />
The metallographic samples for the microstructural examination<br />
were cut from the weld cross-section, then mounted, ground,<br />
polished, <strong>and</strong> etched with a 2% Nital solution. The etched samples<br />
were observed using a light microscope along with a Clemex image<br />
analysis system <strong>and</strong> JEOL JSM-6380 scanning electron microscope<br />
(SEM), equipped with Oxford energy dispersive X-ray spectroscopy<br />
(EDS) <strong>and</strong> three-dimensional fractographic analysis capacity. Vickers<br />
microhardness was measured on the polished samples across<br />
the weld using a computerized microhardness tester at a load <strong>of</strong><br />
200 g <strong>and</strong> a dwell time <strong>of</strong> 15 s. All indentations were adequately<br />
spaced to prevent any potential effect <strong>of</strong> strain fields caused by<br />
adjacent indentations. To ensure the validity <strong>of</strong> test results, two<br />
calibration tests were done using a st<strong>and</strong>ard reference test block<br />
before the microhardness tests on the <strong>welded</strong> joints.<br />
<strong>Tensile</strong> <strong>and</strong> <strong>fatigue</strong> test samples fol<strong>low</strong>ing ASTM-E8/E8M st<strong>and</strong>ard<br />
[15] were sectioned from the <strong>welded</strong> blanks in the rolling<br />
direction (i.e., perpendicular to the welding direction), the examples<br />
<strong>of</strong> which are indicated by the dashed line in Fig. 1a with the<br />
geometry <strong>and</strong> dimensions <strong>of</strong> the test coupons shown in Fig. 1b<br />
<strong>and</strong> c. The test coupons were machined in such a way that the weld<br />
was positioned at the center <strong>of</strong> gauge length for the single linear<br />
weld (Fig. 1b), or two welds were located within the gauge length<br />
symmetrically with respect to the middle <strong>of</strong> the sample (Fig. 1c).<br />
To identify the effect <strong>of</strong> welding <strong>and</strong> its geometry on the tensile<br />
<strong>and</strong> <strong>fatigue</strong> <strong>properties</strong>, coupons <strong>of</strong> the BM were also prepared<br />
<strong>and</strong> tested. <strong>Tensile</strong> tests were conducted using a fully computerized<br />
United tensile testing machine at room temperature <strong>and</strong> with<br />
a range <strong>of</strong> strain rates from 1 10 5 to 1 10 2 s 1 . An extensometer<br />
with a gauge length <strong>of</strong> 50 mm <strong>and</strong> a strain limit <strong>of</strong> 20% was<br />
used to measure the strain during the tensile tests. Load control <strong>fatigue</strong><br />
tests were performed in accordance with ASTM E466 [16] on<br />
a fully computerized Instron 8801 servo-hydraulic testing system.<br />
To avoid potential buckling <strong>of</strong> the test specimens, tension–tension<br />
cyclic loading at a stress ratio <strong>of</strong> R = 0.1 was applied at a frequency<br />
Welding lines<br />
Welding lines<br />
HSLA<br />
HSLA<br />
HSLA<br />
HS HSLA<br />
DP<br />
DP<br />
DP DP DP<br />
Steel<br />
Steel<br />
Steel<br />
LA Steel<br />
Steel<br />
Steel<br />
Steel<br />
Steel<br />
Single linear weld Two linear welds Single linear weld Two linear welds<br />
(a)<br />
Rolling direction<br />
(b)<br />
(c)<br />
Name G W R L A B C<br />
Dimension (mm) 50 12.5 12.5 200 57 50 20<br />
(d)<br />
Fig. 1. Geometry <strong>and</strong> dimensions <strong>of</strong> work pieces <strong>and</strong> test specimens, (a) work pieces to be <strong>welded</strong> using a 6 kW <strong>fiber</strong> <strong>laser</strong>, (b) tensile <strong>and</strong> <strong>fatigue</strong> test samples sectioned from<br />
the single linear weld indicated by the short-dashed lines at (a), (c) tensile <strong>and</strong> <strong>fatigue</strong> test samples cut from the multiple linear weld indicated by the long-dashed lines at (a),<br />
<strong>and</strong> (d) a table <strong>of</strong> the specimen dimensions.
W. Xu et al. / Materials <strong>and</strong> Design 43 (2013) 373–383 375<br />
(a)<br />
(b)<br />
c<br />
e<br />
g<br />
d<br />
f<br />
h<br />
500 µm 500 µm<br />
(c)<br />
(d)<br />
150 µm 150 µm<br />
(e)<br />
(f)<br />
150 µm<br />
150 µm<br />
(g)<br />
(h)<br />
150 µm<br />
150 µm<br />
Fig. 2. Micrographs showing the microstructure change <strong>of</strong> a FLW HSLA <strong>and</strong> a FLW DP980 steel joint, (a) overall view <strong>of</strong> the cross section <strong>of</strong> the FLW HSLA steel joint, (b)<br />
overall view <strong>of</strong> the cross section <strong>of</strong> the FLW DP980 steel joint, (c) HSLA BM, (d) DP980 BM, (e) HAZ <strong>of</strong> the FLW HSLA steel joint, (f) HAZ <strong>of</strong> the FLW DP980 steel joint, (g) FZ <strong>of</strong><br />
the FLW HSLA steel joint <strong>and</strong> (h) FZ <strong>of</strong> the FLW DP980 steel joint.<br />
Table 2<br />
Comparison <strong>of</strong> the average width <strong>of</strong> HAZ <strong>and</strong> FZ in the <strong>welded</strong> joints made with different types <strong>of</strong> <strong>laser</strong> welding.<br />
Welding type<br />
Power<br />
(kW)<br />
Welding speed<br />
(m/min)<br />
Spot size<br />
(mm)<br />
Welding<br />
materials<br />
Thickness <strong>of</strong><br />
workpiece (mm)<br />
Average width <strong>of</strong><br />
HAZ (lm)<br />
Average width <strong>of</strong><br />
FZ (lm)<br />
DLW 4 1.6 12 ⁄ 0.5 DP980 steel 1.2 4000 3000 [11]<br />
Nd:YAG Laser 3 3 0.6 DP980 steel 1.17 1000 750 [10]<br />
welding<br />
CO 2 <strong>laser</strong> 6 6 – DP780 steel 1.8 1000 1000 [12]<br />
welding<br />
FLW 6 16 0.6 DP980 steel 1.2 250 450 Present<br />
study<br />
Reference
376 W. Xu et al. / Materials <strong>and</strong> Design 43 (2013) 373–383<br />
<strong>of</strong> 50 Hz <strong>and</strong> sinusoidal waveform. In both tensile <strong>and</strong> <strong>fatigue</strong> tests,<br />
at least two specimens were tested at each strain rate or each<br />
cyclic stress amplitude. The <strong>fatigue</strong> fracture surfaces were examined<br />
via a scanning electron microscope (SEM) to identify the initiation<br />
sites <strong>and</strong> propagation mechanism <strong>of</strong> the <strong>fatigue</strong> crack.<br />
3. Results <strong>and</strong> discussion<br />
3.1. Microstructure evolution<br />
The representative weld cross-section pr<strong>of</strong>iles <strong>and</strong> microstructure<br />
at different weld zones <strong>of</strong> the HSLA <strong>and</strong> DP980 joints made<br />
with the FLW are shown in Fig. 2. The cross-section <strong>of</strong> the welds<br />
showed a significant microstructural change in both HSLA <strong>and</strong><br />
DP980 joints, indicating the formation <strong>of</strong> the middle FZ <strong>and</strong> two<br />
adjacent HAZs which coexisted with the unaffected BM at both left<br />
<strong>and</strong> right ends as indicated in Fig. 2a <strong>and</strong> b. It is seen that the width<br />
<strong>of</strong> HAZ <strong>and</strong> FZ in both HSLA <strong>and</strong> DP980 joints was similar with an<br />
average <strong>of</strong> 200–300 lm, <strong>and</strong> 400–500 lm, respectively. The HAZ<br />
<strong>and</strong> FZ were narrower when compared to those formed in other <strong>laser</strong><br />
welding (Table 2). For instance, a 4 kW DLW <strong>of</strong> 1.2 mm thick<br />
DP980 steel sheets resulted in 4 mm wide HAZ <strong>and</strong> 3 mm wide<br />
FZ [11]. This was attributed to considerably larger <strong>laser</strong> beam size<br />
(6 mm 2 for the DLW versus 0.28 mm 2 for the FLW), where the <strong>low</strong>er<br />
energy density (less than 10 6 W/cm 2 ) produced conduction<br />
mode. In contrast, in the FLW process the <strong>high</strong>er energy density<br />
which caused keyhole mode made welding more efficient compared<br />
with DLW [10,17,18], <strong>and</strong> the FLW could be operated at<br />
<strong>high</strong>er welding speed with a narrower weld (Table 2).<br />
HSLA BM consisted <strong>of</strong> fine grained ferrite matrix with a uniform<br />
dispersion <strong>of</strong> fine alloy carbides (Fig. 2c), whereas martensite<br />
phase embedded in a continuous ferrite matrix was seen in<br />
DP980 BM (Fig. 2d). The volume fraction <strong>of</strong> martensite in the<br />
DP980 BM was estimated to be 56% using image analysis in<br />
SEM, which matched closely with the earlier studies on the<br />
DP980 steel <strong>of</strong> similar chemistry [13,19–21]. Fig. 2e <strong>and</strong> f showed<br />
the microstructural changes <strong>of</strong> HAZ in both HSLA <strong>and</strong> DP980 joints,<br />
where the typical micrographs were taken from the HAZ<br />
(a)<br />
positioned at the left side <strong>of</strong> FZ corresponding to ‘‘e’’ in Fig. 2a<br />
<strong>and</strong> ‘‘f’’ in Fig. 2b. The details <strong>of</strong> HAZ microstructure will be discussed<br />
in the later section. The microstructure in the FZ depended<br />
heavily on the heat input <strong>and</strong> cooling rate. Based on a continuouscooling<br />
transformation (CCT) diagram <strong>of</strong> weld metal <strong>of</strong> <strong>low</strong> carbon<br />
steel, the FZ microstructure would contain a combination <strong>of</strong> grain<br />
boundary ferrite, side-plate ferrite, accicular ferrite, bainite <strong>and</strong><br />
martensite [22]. However, Fig. 2g <strong>and</strong> h clearly showed that fully<br />
martensitic structure formed in the FZ <strong>of</strong> both HSLA <strong>and</strong> DP980<br />
joints after the present FLW. The reason for this, based on the<br />
CCT diagram <strong>of</strong> weld metal <strong>of</strong> <strong>low</strong> carbon steel [22], was believed<br />
to be associated with the much <strong>high</strong>er power density <strong>and</strong> rapid<br />
cooling rate in the FZ during FLW, where the cooling curve being<br />
steeper enough missed to touch the nose <strong>of</strong> the C-curves, thus<br />
resulting in only martensite transformation. Sreenivasan et al.<br />
[10] reported a similar result for a 3 kW Nd:YAG <strong>laser</strong> welding applied<br />
on the DP980 steel sheets with a thickness <strong>of</strong> 1.17 mm. It<br />
should be noted that similar microstructural change was observed<br />
in the case <strong>of</strong> the multiple linear welds except the presence <strong>of</strong><br />
twice microstructural heterogeneities, i.e., FZ <strong>and</strong> HAZ, because<br />
<strong>of</strong> doubling the linear welds. Thus, their microstructure was not<br />
presented here.<br />
SEM images <strong>of</strong> the inter-critical HAZ (experiencing temperatures<br />
between Ac 1 <strong>and</strong> Ac 3 lines) with adjacent BM <strong>and</strong> upper-critical HAZ<br />
(region above Ac 3 line) <strong>of</strong> HSLA joint are shown in Fig. 3. The HSLA<br />
BM microstructure (Fig. 3b was confirmed to consist <strong>of</strong> ferrite grains<br />
with finely dispersed alloy carbide particles as observed in optical<br />
image (Fig. 2c). Fig. 3c represented the microstructure containing<br />
some martensite which formed as a product <strong>of</strong> fast cooling from<br />
the inter-critical region, where the partially formed austenite (i.e.,<br />
co-existence <strong>of</strong> ferrite <strong>and</strong> austenite) transformed to martensite<br />
during fast cooling down to room temperature in the FLW process.<br />
However, it should be pointed out that the HAZ <strong>of</strong> HSLA joint did not<br />
involve any microstructural changes at the sub-critical region (be<strong>low</strong><br />
Ac 1 line) due to the absence <strong>of</strong> martensite in its BM.<br />
On the other h<strong>and</strong>, for DP980 steel after FLW a tempered martensite<br />
area (or sub-critical region be<strong>low</strong> Ac 1 line) located in-between<br />
the BM <strong>and</strong> inter-critical zone (between Ac 1 <strong>and</strong> Ac 3<br />
(b)<br />
b<br />
c<br />
F<br />
BM<br />
Inter-critical HAZ<br />
Uppercritical<br />
HAZ<br />
F<br />
(c)<br />
F<br />
M<br />
Fig. 3. SEM micrographs showing the HAZ microstructure change in the FLW <strong>of</strong> HSLA steel, (a) overall view <strong>of</strong> inter-critical HAZ with adjacent BM <strong>and</strong> upper-critical HAZ, (b)<br />
HSLA BM with ferrite <strong>and</strong> fine alloy carbides, <strong>and</strong> (c) inter-critical HAZ showing newly formed martensite phase (M: martensite, <strong>and</strong> F: ferrite).
W. Xu et al. / Materials <strong>and</strong> Design 43 (2013) 373–383 377<br />
(a)<br />
b<br />
c<br />
(b)<br />
F<br />
M<br />
BM<br />
Tempered<br />
area (subcritical<br />
HAZ)<br />
Intercritical<br />
HAZ<br />
Uppercritical<br />
HAZ<br />
(c)<br />
F<br />
TM<br />
Fig. 4. SEM micrographs showing the microstructural change in a DP980 steel joint, (a) overall view <strong>of</strong> tempered martensite area in the HAZ with adjacent BM, inter-critical<br />
HAZ, <strong>and</strong> upper-critical HAZ, (b) DP980 BM showing ferrite <strong>and</strong> martensite, <strong>and</strong> (c) tempered martensite region in the HAZ (M: martensite, TM: tempered martensite, <strong>and</strong> F:<br />
ferrite).<br />
lines), along with an upper-critical HAZ (above Ac 3 line) could be<br />
seen in Fig. 4a. While the microstructure in the inter-critical zone<br />
was somewhat similar to that in the BM with ferrite plus martensite<br />
isl<strong>and</strong>s, in the tempered martensite area some extent <strong>of</strong> decomposition<br />
<strong>of</strong> the martensite pre-existed in the BM (Fig. 4b) was<br />
clearly visible at a <strong>high</strong>er magnification as shown in Fig. 4c. Tempered<br />
martensite <strong>and</strong> the resultant s<strong>of</strong>tening zone in the <strong>laser</strong><br />
welding <strong>of</strong> DP980 steel have been reported in many earlier studies<br />
[11,13,14,23]. In general, the temperature undergone in the s<strong>of</strong>tening<br />
region during welding was close to or be<strong>low</strong> the critical temperature<br />
(Ac 1 at which austenite begins to form during heating)<br />
which led to <strong>high</strong> temperature tempering <strong>of</strong> martensite phase in<br />
the BM <strong>of</strong> DP steels [11,13,14,20,23–25]. The microstructure <strong>of</strong><br />
tempered martensite in DP980 welds has been reported to consist<br />
typically <strong>of</strong> uniformly distributed fine precipitated cementite<br />
(Fe 3 C) particles embedded within a continuous ferrite matrix<br />
[20,26–28]. As seen from Fig. 4c, it was clear that only partial tempering<br />
<strong>of</strong> martensite occurred in the HAZ <strong>of</strong> DP980 in the FLW<br />
process.<br />
Vickers hardness, HV<br />
600<br />
500 BM HAZ FZ HAZ BM<br />
400<br />
FLW HSLA-S<br />
FLW DP980-S<br />
300<br />
S<strong>of</strong>t zone S<strong>of</strong>t zone<br />
200<br />
100 BM HAZ FZ HAZ BM<br />
0<br />
-6 -4 -2 0 2 4 6<br />
Distance from weld centreline, mm<br />
Fig. 5. Representative hardness pr<strong>of</strong>iles across the HSLA-S <strong>and</strong> DP980-S joints.<br />
3.2. Hardness pr<strong>of</strong>iles<br />
Vickers microhardness pr<strong>of</strong>iles across the welds <strong>of</strong> the FLW<br />
HSLA single linear joint (FLW HSLA-S) <strong>and</strong> the FLW DP980 single<br />
linear joint (FLW DP980-S) (as indicated by the red 1 dashed lines<br />
in Fig 2a <strong>and</strong> b) are showed in Fig. 5. The average hardness value<br />
<strong>of</strong> the HSLA <strong>and</strong> DP980 steel BM was measured to be about<br />
170 HV <strong>and</strong> 340 HV, respectively. The <strong>high</strong>er hardness in the<br />
DP980 BM was attributed to the presence <strong>of</strong> <strong>high</strong>er fraction <strong>of</strong> hard<br />
martensite phase (56%, Fig. 2d <strong>and</strong> Fig. 4b), while the HSLA BM<br />
did not contain any martensite (Fig. 2c <strong>and</strong> Fig. 3b). In accordance<br />
with the microstructure (Fig. 2g <strong>and</strong> h) where fully martensitic<br />
structure emerged, the hardness <strong>of</strong> the FZ became much <strong>high</strong>er<br />
1 For interpretation <strong>of</strong> color in Figs. 2-11, the reader is referred to the web version<br />
<strong>of</strong> this article.<br />
<strong>and</strong> showed a decrease from FZ to the HAZ. Interestingly, the hardness<br />
curve though the whole HAZ was seen to merge smoothly into<br />
the unaffected BM in the HSLA joint due to a decrease in the fraction<br />
<strong>of</strong> martensite in the HAZ which formed as a solid-state phase<br />
transformation from austenite in the inter-critical region. On the<br />
other h<strong>and</strong>, the hardness pattern in the DP980 joint exhibited<br />
two s<strong>of</strong>t zone ‘‘valleys’’ beside the center FZ, where the hardness<br />
locally dropped significantly to an average <strong>of</strong> 280 HV which was<br />
well be<strong>low</strong> the DP980 BM hardness (350 HV). This was attributed<br />
to the occurrence <strong>of</strong> martensite tempering (HAZ-s<strong>of</strong>tened region)<br />
as discussed above <strong>and</strong> also illustrated in Fig. 4. The narrow FZ<br />
with fully martensitic microstructure in both HSLA <strong>and</strong> DP980<br />
joints after FLW were characterized by a very <strong>high</strong> hardness<br />
(Fig. 5). However, the DP980 joint had a further <strong>high</strong>er hardness<br />
averaging 450 HV compared with the average FZ hardness<br />
(360 HV) in the HSLA joint. This was mainly due to the fact that
378 W. Xu et al. / Materials <strong>and</strong> Design 43 (2013) 373–383<br />
Vickers hardness, HV<br />
600<br />
500<br />
400<br />
300<br />
200<br />
100<br />
S<strong>of</strong>t<br />
zone<br />
S<strong>of</strong>t<br />
zone<br />
FLW HSLA-M<br />
FLW DP980-M<br />
S<strong>of</strong>t<br />
zone<br />
0<br />
-4 0 4 8 12 16 20 24<br />
Distance from left weld line, mm<br />
S<strong>of</strong>t<br />
zone<br />
BM HAZ FZ HAZ BM HAZ FZ HAZ BM<br />
observed to display a ‘‘suspension bridge’’-like or ‘‘cable-stayed<br />
bridge’’-like hardness pr<strong>of</strong>ile with two FZ hardness representing<br />
the ‘‘pylons or towers’’. It is seen that each weld bead <strong>of</strong> the multiple<br />
linear joints (Fig. 6) had a similar hardness values when compared<br />
it to the single linear joints (Fig. 5), which was a result <strong>of</strong><br />
similar microstructures as noted in the previous section. The distinction<br />
between the FLW HSLA-M <strong>and</strong> FLW DP980-M joints was<br />
that the hardness pr<strong>of</strong>ile <strong>of</strong> the former was smoother without ‘‘valleys’’<br />
due to the absence <strong>of</strong> HAZ-s<strong>of</strong>tening (martensite tempering).<br />
On the contrary, in the FLW DP980-M joint each FZ hardness ‘‘pylon’’<br />
was accompanied by two s<strong>of</strong>t zone ‘‘valleys’’ positioned at its<br />
two sides (Fig. 6). It would be <strong>of</strong> interest to see if doubling the<br />
number <strong>of</strong> s<strong>of</strong>t zones would affect tensile <strong>strength</strong> <strong>and</strong> <strong>fatigue</strong> life<br />
<strong>of</strong> the HSLA <strong>and</strong> DP980 joints.<br />
3.3. <strong>Tensile</strong> <strong>properties</strong><br />
Fig. 6. Typical hardness pr<strong>of</strong>iles across the HSLA-M <strong>and</strong> DP980-M joints.<br />
Engineering stress, MPa<br />
1200<br />
900<br />
600<br />
300<br />
0<br />
0 5 10 15 20<br />
Engineering strain, %<br />
HSLA BM<br />
FLW HSLA-S<br />
FLW HSLA-M<br />
DP980 BM<br />
FLW DP980-S<br />
FLW DP980-M<br />
Fig. 7. Representative engineering stress versus engineering strain curves <strong>of</strong> HSLA<br />
BM, FLW HSLA-S joint, FLW HSLA-M joint, DP980 BM, FLW DP980-S joint, FLW<br />
DP980-M joint tested at a strain rate <strong>of</strong> 1 10 3 s 1 .<br />
the martensite in the FZ <strong>of</strong> DP980 joint contained a <strong>high</strong>er amount<br />
<strong>of</strong> carbon, as seen in Table 1 where nearly twice carbon content<br />
was present in the DP980 steel compared with the HSLA steel.<br />
Additionally, nearly doubled manganese content in the DP980 steel<br />
(Table 1) also contributed to the <strong>high</strong>er hardness [29]. Calcagnotto<br />
et al. [30] noted that Mn was able to increase hardneability<br />
substantially just like carbon content in austenite on grain size stability<br />
<strong>and</strong> hardneability in ultrafine-grained ferrite/martensite<br />
dual-phase steel.<br />
Fig. 6 shows the microhardness pr<strong>of</strong>iles across the multiple linear<br />
weld <strong>of</strong> FLW HSLA (FLW HSLA-M) <strong>and</strong> FLW DP980 (FLW<br />
DP980-M) joints. Both FLW HSLA-M <strong>and</strong> FLW DP980-M joints were<br />
The representative engineering stress versus engineering strain<br />
curves, determined at a strain rate <strong>of</strong> 1 10 3 s 1 , <strong>of</strong> the BM, FLW<br />
single <strong>and</strong> multiple linear joints <strong>of</strong> both HSLA <strong>and</strong> DP980 steels are<br />
illustrated in Fig. 7. In the HSLA joints, both the strain to failure <strong>and</strong><br />
the yield <strong>strength</strong> (YS) after welding were observed to be very<br />
close to those <strong>of</strong> the BM. The evaluated tensile test results are<br />
listed in Table 3. The YS <strong>and</strong> ultimate tensile <strong>strength</strong> (UTS) were<br />
obtained to be 400 MPa <strong>and</strong> 515 MPa, respectively, for the FLW<br />
HSLA-S joint <strong>and</strong> 398 MPa <strong>and</strong> 523 MPa, respectively, for the<br />
FLW HSLA-M joint. These values were fairly close to those <strong>of</strong> the<br />
BM which had a YS <strong>of</strong> 455 MPa <strong>and</strong> UTS <strong>of</strong> 546 MPa (Table 3, <strong>and</strong><br />
Fig. 7). These tensile results suggested that the present FLW process,<br />
regardless <strong>of</strong> single or multiple linear welding had little influence<br />
on the tensile <strong>properties</strong> <strong>of</strong> the HSLA steel. Obviously, this was<br />
because the FLW <strong>of</strong> HSLA steel did not create any s<strong>of</strong>t zone (Fig. 5<br />
<strong>and</strong> 6) in the HAZ that was detrimental to the mechanical <strong>properties</strong><br />
<strong>of</strong> the welds. This was also corroborated by the tensile failure<br />
<strong>of</strong> FLW HSLA joints that consistently occurred in the BM. This finding<br />
was also in agreement with the earlier observations on the tensile<br />
failure locations by Xia et al. in a DLW HSLA joint [11], Farabi<br />
et al. [31] in a diode <strong>laser</strong> <strong>welded</strong> DP600 steel joint, <strong>and</strong> P<strong>and</strong>a et al.<br />
[32] in a diode <strong>laser</strong> <strong>welded</strong> HSLA/DP980 dissimilar joint with a<br />
thickness <strong>of</strong> 1.14 mm <strong>and</strong> 1.17 mm, respectively.<br />
In contrast to the HSLA welds, the DP980 joints showed a <strong>low</strong>er<br />
strain to failure (Fig. 7 <strong>and</strong> Table 3). However, the YS <strong>and</strong> UTS were<br />
basically unaffected by the FLW. As seen in Table 2, the <strong>strength</strong> <strong>of</strong><br />
FLW DP980-S joints (YS = 719 MPa; UTS = 1048 MPa) <strong>and</strong> FLW<br />
DP980-M joints (YS = 717 MPa, UTS = 1066 MPa) was very close<br />
to those <strong>of</strong> the BM (YS = 720 MPa; UTS = 1095 MPa). The results<br />
indicated that the presence <strong>of</strong> narrow s<strong>of</strong>t zone in the FLW<br />
DP980 welds was not detrimental to the tensile <strong>strength</strong>, <strong>and</strong> the<br />
<strong>high</strong> welding speed could greatly reduce the harmful influence <strong>of</strong><br />
the s<strong>of</strong>t zone by narrowing it down significantly.<br />
It was also <strong>of</strong> interest to observe that a joint efficiency (i.e., a ratio<br />
<strong>of</strong> the UTS <strong>of</strong> <strong>welded</strong> joints to the UTS <strong>of</strong> the corresponding BM<br />
[33–35]) <strong>of</strong> both FLW HSLA joints <strong>and</strong> FLW DP980 joints reached as<br />
<strong>high</strong> as about 94–96% <strong>and</strong> 96–97% (Table 3), respectively. These<br />
Table 3<br />
<strong>Tensile</strong> <strong>properties</strong> <strong>and</strong> <strong>fatigue</strong> parameters r 0 f<br />
<strong>and</strong> b for HSLA BM, FLW HSLA-S joint, FLW HSLA-M joint, DP980 BM, FLW DP980-S joint, <strong>and</strong> FLW DP980-M joint.<br />
Welding type<br />
Yield <strong>strength</strong><br />
(MPa)<br />
Ultimate tensile <strong>strength</strong><br />
(MPa)<br />
Elongation<br />
(%)<br />
Joint efficiency<br />
(%)<br />
Fatigue limit<br />
(MPa)<br />
Fatigue<br />
ratio<br />
r 0 f<br />
(MPa)<br />
b<br />
HSLA BM 455 546 24.7 – 200 0.366 302 0.023<br />
FLW HSLA-S 400 515 22.7 94.3 125 0.243 328 0.041<br />
FLW HSLA-M 398 523 20.7 95.8 125 0.239 343 0.048<br />
DP980 BM 720 1095 14.2 – 250 0.228 1019 0.098<br />
FLW DP980-S 719 1048 4.6 95.7 150 0.143 1169 0.132<br />
FLW DP980-M 717 1066 5.1 97.3 150 0.141 1984 0.185
W. Xu et al. / Materials <strong>and</strong> Design 43 (2013) 373–383 379<br />
results were similar to those <strong>of</strong> Nd:YAG <strong>laser</strong> welding (94.1% for<br />
DP980 steel joints) [12] <strong>and</strong> CO 2 <strong>laser</strong> welding (96.9% for CP1180<br />
steel joints [12]). Indeed, if one considers the ratio <strong>of</strong> the YS <strong>of</strong><br />
<strong>welded</strong> joints to the YS <strong>of</strong> the corresponding BM, a nearly 100%<br />
YS joint efficiency would be achieved in both FLW DP980-S <strong>and</strong><br />
FLW DP980-M joints. From these tensile results it can be concluded<br />
that FLW is an able <strong>and</strong> viable joining process that gave rise to<br />
superior tensile performance approaching that <strong>of</strong> BMs for both<br />
HSLA <strong>and</strong> DP980 steels, while the FLW DP980 joints had a much<br />
<strong>high</strong>er tensile <strong>strength</strong> than the FLW HSLA joints.<br />
However, it should be noted that, while the percent elongation<br />
or ductility <strong>of</strong> both FLW HSLA-S <strong>and</strong> HSLA-M joints (21–23%)<br />
came fairly close to that <strong>of</strong> the HSLA BM (25%) as shown in Table<br />
3, the ductility <strong>of</strong> the FLW DP980-S joints (4.6%) <strong>and</strong> the FLW<br />
DP980-M joints (5.1%) was noticeably <strong>low</strong>er than that <strong>of</strong> the BM<br />
(14.2%) since all the tensile test samples failed consistently in the<br />
s<strong>of</strong>t zone. The <strong>low</strong>er ductility in the DP980 welds was due to the<br />
fact that yielding occurred first in one <strong>of</strong> the s<strong>of</strong>t zones <strong>and</strong> the<br />
subsequent plastic deformation was predominantly concentrated<br />
there leading to necking <strong>and</strong> premature failure at that location,<br />
which resulted in a reduction in the overall specimen elongation.<br />
Similar results have also been reported in the study <strong>of</strong> formability<br />
<strong>of</strong> DLW <strong>of</strong> DP980 [11,32]. It was also due to the premature failure<br />
induced by the s<strong>of</strong>t zone (i.e., <strong>low</strong>er ductility) that the UTS <strong>of</strong> FLW<br />
joints became <strong>low</strong>er (Fig. 7), although a <strong>high</strong> joint efficiency <strong>of</strong><br />
about 96–97% was achieved in the FLW DP980 joints.<br />
3.4. Fatigue <strong>strength</strong> <strong>and</strong> failure mode<br />
Fatigue test results <strong>of</strong> the BM, FLW single linear <strong>and</strong> multiple<br />
linear joints <strong>of</strong> HSLA tested at R = 0.1, 50 Hz, <strong>and</strong> room temperature<br />
(RT) are plotted in the Fig. 8a. Both FLW HSLA-S <strong>and</strong> HSLA-M joints<br />
exhibited an equivalent <strong>fatigue</strong> life within the experimental scatter.<br />
At a stress amplitude <strong>of</strong> above 250 MPa the <strong>fatigue</strong> life <strong>of</strong> the<br />
joints was only slightly <strong>low</strong>er than that <strong>of</strong> the BM, whereas the <strong>fatigue</strong><br />
<strong>strength</strong> became <strong>low</strong>er <strong>and</strong> more scattered at <strong>low</strong>er stress<br />
amplitudes (
380 W. Xu et al. / Materials <strong>and</strong> Design 43 (2013) 373–383<br />
stress amplitudes increased with increasing number <strong>of</strong> s<strong>of</strong>t zones<br />
in the <strong>high</strong>er grade <strong>of</strong> DP980.<br />
To make a better comparison, <strong>fatigue</strong> test results <strong>of</strong> the FLW<br />
HSLA-S joints, the FLW HSLA-M joints, the FLW DP980-S joints,<br />
<strong>and</strong> the FLW DP980-M joints were plotted in Fig. 8c. The FLW<br />
DP980 joints were observed to have a much <strong>high</strong>er <strong>fatigue</strong> <strong>strength</strong><br />
than that <strong>of</strong> the FLW HSLA joints at a stress amplitude above<br />
250 MPa, but the <strong>fatigue</strong> <strong>strength</strong> was equivalent be<strong>low</strong> a stress<br />
amplitude <strong>of</strong> 250 MPa for both DP980 <strong>and</strong> HSLA joints due to the<br />
fact that the <strong>high</strong>er grade/<strong>strength</strong> <strong>of</strong> DP980 was more sensitive<br />
to the presence <strong>of</strong> severe weld concavity (Fig. 2b) especially at<br />
the <strong>low</strong>er level <strong>of</strong> cyclic load than at the <strong>high</strong>er level <strong>of</strong> cyclic load.<br />
Fatigue limit <strong>and</strong> <strong>fatigue</strong> ratio are tabulated in Table 3. The decrease<br />
in <strong>fatigue</strong> limit for both FLW HSLA <strong>and</strong> FLW DP980 joints,<br />
with respect to their BM, were similar at 37.5% <strong>and</strong> 40%, respectively,<br />
while the <strong>fatigue</strong> limit <strong>of</strong> the FLW DP980 joints were 20%<br />
<strong>high</strong>er than that <strong>of</strong> the FLW HSLA joints. The <strong>fatigue</strong> ratio (i.e., a ratio<br />
<strong>of</strong> <strong>fatigue</strong> limit to the UTS) <strong>of</strong> the FLW HSLA-S joints <strong>and</strong> the<br />
FLW DP980-S joints were 0.243 <strong>and</strong> 0.143, respectively. The <strong>low</strong>er<br />
<strong>fatigue</strong> ratio <strong>of</strong> FLW DP980-S joints was related to its much <strong>high</strong>er<br />
(or twice) UTS (Table 3). These results suggested that while the<br />
tensile <strong>strength</strong> <strong>of</strong> the FLW DP980 joints was much <strong>high</strong>er than<br />
that <strong>of</strong> the FLW HSLA joints, the <strong>fatigue</strong> <strong>strength</strong> was very close<br />
for both joints at a <strong>low</strong>er level <strong>of</strong> stress amplitudes (
W. Xu et al. / Materials <strong>and</strong> Design 43 (2013) 373–383 381<br />
in both joints was similar to a notched specimen studied by Chapetti<br />
et al. [38] where the stress concentration was generated,<br />
<strong>and</strong> consequently a decreased <strong>fatigue</strong> limit or <strong>fatigue</strong> life at the<br />
<strong>low</strong> level <strong>of</strong> stress amplitudes was observed. Further studies are<br />
needed to eliminate the weld concavity <strong>and</strong> improve the <strong>fatigue</strong><br />
resistance.<br />
The obtained <strong>fatigue</strong> data plotted in Fig. 8 could be fitted using<br />
the fol<strong>low</strong>ing Basquin type equation,<br />
r a ¼ r 0 f ð2NÞb<br />
where r a is the stress amplitude, r 0 f<br />
is the <strong>fatigue</strong> <strong>strength</strong> coefficient<br />
defined by the stress intercept at 2N =1,N is the number <strong>of</strong><br />
cycle to failure, <strong>and</strong> b is the <strong>fatigue</strong> <strong>strength</strong> exponent. The obtained<br />
values <strong>of</strong> r 0 f<br />
<strong>and</strong> b <strong>of</strong> the BM, the FLW single linear joint <strong>and</strong> the FLW<br />
multiple linear joints for both HSLA <strong>and</strong> DP980 steels tested at<br />
R = 0.1, 50 Hz <strong>and</strong> room temperature are given in Table 3. Apparently,<br />
the <strong>fatigue</strong> life at a given stress amplitude was dependent<br />
on both <strong>fatigue</strong> <strong>strength</strong> coefficient r 0 f<br />
<strong>and</strong> <strong>fatigue</strong> <strong>strength</strong> exponent<br />
b with b being a predominant factor. The smaller the absolute<br />
value <strong>of</strong> b, the longer the <strong>fatigue</strong> life. It is seen from Table 3 that<br />
after FSW for both HSLA <strong>and</strong> DP980 steels the absolute values <strong>of</strong><br />
b became larger, giving rise to a <strong>low</strong>er <strong>fatigue</strong> life after welding.<br />
3.5. Failure location <strong>and</strong> fractography<br />
The typical <strong>fatigue</strong> failure locations for the HSLA-S joints, the<br />
HSLA-M joints, the DP980-S joints <strong>and</strong> the DP980-M joints are<br />
shown in Fig. 9. It is seen that the <strong>fatigue</strong> failure for the FLW HSLA<br />
joints occurred in the BM with large deformation at a stress amplitude<br />
above 250 MPa (Fig. 9a <strong>and</strong> b), whereas it occurred at HAZ be<strong>low</strong><br />
a stress amplitude <strong>of</strong> 250 MPa (Fig. 9c <strong>and</strong> d), irrespective <strong>of</strong><br />
the single or multiple linear welding. It was believed that while<br />
ð1Þ<br />
the weld concavity in the FLW HSLA joints (Fig. 2a) did not significantly<br />
influence the tensile <strong>strength</strong> (Fig. 7 <strong>and</strong> Table 3), the <strong>fatigue</strong><br />
<strong>strength</strong> at the condition <strong>of</strong> stress amplitude be<strong>low</strong> 250 MPa<br />
was affected adversely. In contrast, the <strong>fatigue</strong> failure in the FLW<br />
DP980 joints always initiated in weld concavity (Fig. 9e <strong>and</strong> f)<br />
due to the presence <strong>of</strong> stress concentration, <strong>and</strong> then propagated<br />
in the s<strong>of</strong>t zone (Fig. 5 <strong>and</strong> 6). The failure location <strong>of</strong> the FLW joints<br />
was similar to that <strong>of</strong> the DLW DP980 joints in which all the <strong>fatigue</strong><br />
failure occurred at the s<strong>of</strong>t zones as reported in our earlier study<br />
[39].<br />
SEM images <strong>of</strong> a fracture surface <strong>of</strong> a FLW HSLA-S joint tested at<br />
an applied stress amplitude <strong>of</strong> 200 MPa is shown in Fig. 10. The<br />
cross section in Fig 10a indicated that the <strong>fatigue</strong> test sample experienced<br />
a large deformation before the final fast crack propagation<br />
because <strong>of</strong> the superior ductility <strong>of</strong> HSLA steel in comparison with<br />
DP980 steel (Table 3). It was also confirmed from Fig. 10a <strong>and</strong> b<br />
that crack initiation occurred from the weld concavity (Fig. 2a),<br />
corresponding well to the top view <strong>of</strong> the failed samples (Fig. 9c<br />
<strong>and</strong> d). A magnified SEM image near the crack initiation site is<br />
shown in Fig. 10b, where the crack initiated from the weld concavity<br />
could clearly be seen when the applied stress amplitude was<br />
be<strong>low</strong> 250 MPa. Fatigue crack propagation was mainly characterized<br />
by the formation <strong>of</strong> fairly typical <strong>fatigue</strong> striations in conjunction<br />
with secondary cracks, as shown in Fig. 10c, while the final fast<br />
propagation area consisted <strong>of</strong> distinctive dimples (Fig. 10d). Fatigue<br />
striations basically occurred through a repeated plastic blunting–sharpening<br />
process via the slip <strong>of</strong> dislocations in the plastic<br />
zone in front <strong>of</strong> the <strong>fatigue</strong> crack tip [40].<br />
SEM images <strong>of</strong> a fracture surface <strong>of</strong> a FLW DP980-S joint tested<br />
at an applied stress amplitude <strong>of</strong> 200 MPa is shown in Fig. 11. Itis<br />
from Fig. 11a that multiple crack initiation occurred from the weld<br />
concavity where the surface welding defect occurs (Fig. 11b) as<br />
indicated by the arrow in the image. These characteristics <strong>of</strong> the<br />
(a)<br />
b<br />
d<br />
1 mm<br />
(b)<br />
(c)<br />
c<br />
Surface defect<br />
(d)<br />
Fig. 11. Typical SEM images <strong>of</strong> <strong>fatigue</strong> fracture surface <strong>of</strong> a FLW DP980-S joint tested at a stress amplitude 200 MPa, (a) overall view, (b) crack initiation area, (c) crack<br />
propagation area, <strong>and</strong> (d) final fast crack propagation area.
382 W. Xu et al. / Materials <strong>and</strong> Design 43 (2013) 373–383<br />
<strong>fatigue</strong> fractograph were similar to those observed on the <strong>fatigue</strong><br />
fracture surface reported for a FLW AZ31B-H24 Mg alloy [35].<br />
The <strong>fatigue</strong> crack initiation in the present FLW DP980 steel would<br />
primarily be related to the stress concentration caused by the severe<br />
weld concavity (Fig. 2b), in combination with the presence<br />
<strong>of</strong> detrimental s<strong>of</strong>t zones (Fig. 5 <strong>and</strong> 6). Like the HSLA welds, <strong>fatigue</strong><br />
crack propagation in the FLW DP980 joints was also characterized<br />
by the occurrence <strong>of</strong> <strong>fatigue</strong> striations along with secondary<br />
cracks, as shown in Fig. 11c, whereas the final fast propagation area<br />
consisted <strong>of</strong> characteristic dimples as well, as shown in Fig. 11d,<br />
where some remaining inclusions located at the bottom <strong>of</strong> dimples<br />
could be seen.<br />
4. Conclusions<br />
In this study the microstructure, hardness pr<strong>of</strong>ile, tensile <strong>properties</strong>,<br />
<strong>and</strong> <strong>fatigue</strong> performance <strong>of</strong> <strong>high</strong> speed <strong>fiber</strong> <strong>laser</strong> <strong>welded</strong><br />
HSLA <strong>and</strong> DP980 steel joints with single linear <strong>and</strong> multiple linear<br />
welds are evaluated <strong>and</strong> compared. The fol<strong>low</strong>ing conclusions can<br />
be drawn:<br />
(1) Single phase microstructure containing martensite formed<br />
in the FZ <strong>of</strong> all the joints; however, newly formed martensite<br />
<strong>and</strong> partially tempered martensite structure was observed in<br />
the (HAZ) <strong>of</strong> HSLA <strong>and</strong> DP980 steels, respectively.<br />
(2) The HSLA joints exhibited a hardness pr<strong>of</strong>ile with HAZ hardness<br />
merging smoothly into the unaffected BM values, while<br />
the DP980 joints showed a significant drop <strong>of</strong> hardness in<br />
the HAZ. Despite the formation <strong>of</strong> fully martensitic structure<br />
in the FZ, the <strong>high</strong>er carbon content in DP980 steel resulted<br />
in a <strong>high</strong>er FZ hardness (450 HV) compared to HSLA steel<br />
(360 HV). A characteristic ‘‘suspension bridge’’-like<br />
hardness pr<strong>of</strong>ile with the FZ hardness as a ‘‘pylon’’ was<br />
observed.<br />
(3) Both the HSLA <strong>and</strong> DP980 joints showed a superior tensile<br />
<strong>strength</strong> with a joint efficiency reaching about 94–97%.<br />
While the ductility <strong>of</strong> DP980 joints decreased, the ductility<br />
<strong>of</strong> HSLA joints was fairly close to that <strong>of</strong> BM. Despite<br />
the presence <strong>of</strong> the s<strong>of</strong>t zone, the ultimate tensile<br />
<strong>strength</strong> <strong>of</strong> the DP980 joints was twice <strong>high</strong>er than that<br />
<strong>of</strong> the HSLA joints, which corresponded well to the hardness<br />
results.<br />
(4) The DP980 joints had a much longer <strong>fatigue</strong> life compared to<br />
the HSLA joints at a stress amplitude above 250 MPa; however,<br />
be<strong>low</strong> this stress amplitude both the DP980 <strong>and</strong> HSLA<br />
joints showed a similar <strong>fatigue</strong> life.<br />
(5) No significant effect <strong>of</strong> weld geometry on the <strong>fatigue</strong> life <strong>of</strong><br />
HSLA joints was observed. On the other h<strong>and</strong>, the <strong>fatigue</strong> life<br />
for the multiple linear DP980 joints exhibited a larger scatter<br />
<strong>and</strong> also <strong>low</strong>er <strong>fatigue</strong> <strong>strength</strong>, indicating that the probability<br />
<strong>of</strong> dynamic <strong>fatigue</strong> failure at the <strong>low</strong>er level <strong>of</strong> stress<br />
amplitudes increased with increasing number <strong>of</strong> s<strong>of</strong>t zone<br />
<strong>and</strong> weld concavity in the weld.<br />
(6) Above a stress amplitude <strong>of</strong> 250 MPa, <strong>fatigue</strong> failure <strong>of</strong> the<br />
HSLA joints occurred in the BM, signifying a superior structural<br />
integrity <strong>of</strong> the <strong>fiber</strong> <strong>laser</strong> <strong>welded</strong> HSLA joints when<br />
subjected to <strong>high</strong>er levels <strong>of</strong> cyclic loading. When the<br />
applied stress amplitude was be<strong>low</strong> 250 MPa, <strong>fatigue</strong> failure<br />
occurred from the weld concavity.<br />
(7) In the DP980 joints under both <strong>high</strong>er <strong>and</strong> <strong>low</strong>er levels <strong>of</strong><br />
cyclic loading, <strong>fatigue</strong> crack was observed to initiate predominantly<br />
from the weld concavity <strong>and</strong> propagate in the<br />
s<strong>of</strong>t zone. Fatigue crack propagation in both HSLA <strong>and</strong><br />
DP980 joints was characterized by the characteristic <strong>fatigue</strong><br />
striations coupled with secondary cracks.<br />
Acknowledgements<br />
The authors would like to thank the Natural Sciences <strong>and</strong> Engineering<br />
Research Council <strong>of</strong> Canada (NSERC) <strong>and</strong> AUTO21 Network<br />
<strong>of</strong> Centers <strong>of</strong> Excellence for providing financial support. The financial<br />
support from International Zinc Association (IZA) <strong>and</strong> Arcelor-<br />
Mittal D<strong>of</strong>asco is <strong>high</strong>ly acknowledged. One <strong>of</strong> the authors (D.L.<br />
Chen) is grateful for the financial support by the Premier’s Research<br />
Excellence Award (PREA), NSERC-Discovery Accelerator<br />
Supplement (DAS) Award, Canada Foundation for Innovation<br />
(CFI), <strong>and</strong> Ryerson Research Chair (RRC) program. The authors<br />
would like to thank Dr. J. Chen <strong>and</strong> Dr. Y.L. He (CANMET-Materials<br />
Technology Laboratory, Natural Resources Canada, Hamilton, Canada),<br />
Mr. E. Biro (ArcelorMittal Global Research, Hamilton, Canada),<br />
<strong>and</strong> Dr. J. Villafuerte (CenterLine (Windsor) Ltd., Windsor, Canada)<br />
for their support <strong>and</strong> helpful discussion. The assistance <strong>of</strong> Q. Li, A.<br />
Machin, J. Amankrah, <strong>and</strong> R. Churaman in performing the experiments<br />
is gratefully acknowledged.<br />
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